Hot rolled coated steel sheet having high strength, high formability, excellent bake hardenability and method of manufacturing same

ABSTRACT

A hot-rolled coated steel sheet including: in wt %, C: 0.05-0.14%, Si: 0.1-1.0%, Mn: 1.0-2.0%, P: 0.001-0.05%, S: 0.001-0.01%, AI: 0.01-0.1%, Cr: 0.005-1.0%, Ti: 0.005-0.13%, Nb: 0.005-0.03%, N: 0.001-0.01%, Fe residues, and other inevitable impurities; a mixed structure of ferrite and bainite as a main phase; and as a remaining structure, one or more selected from the group consisting of martensite, austenite, and phase martensite (MA), wherein a fraction of the ferrite and bainite is 95-99 area % and Equation 1 is satisfied. [Equation 1] FCO{110}&lt;112&gt;+FCO{112}&lt;111&gt;≥10 where, FCO{110}&lt;112&gt; and FCO{112}&lt;111&gt;, each representing an area fraction occupied by a structure having ac crystal orientation of {110}&lt;112&gt; and {112}&lt;111&gt;.

CROSS-REFERENCE OF RELATED APPLICATIONS

This application is the U.S. National Phase under 35 U.S.C. § 371 ofInternational Patent Application No. PCT/KR2019/008378, filed on Jul. 8,2019, which in turn claims the benefit of Korean Application No.10-2018-0081086, filed on Jul. 12, 2018, the entire disclosures of whichapplications are incorporated by reference herein.

TECHNICAL FIELD

The present invention relates to a hot-rolled coated steel sheet havinghigh strength, high formability, and excellent bake hardenability, and amethod of manufacturing the same, and more particularly, to a hot-rolledcoated steel sheet that can be preferably applied to brackets,reinforcing materials, connecting materials, and the like, of automobilechassis parts, and a method of manufacturing the same.

BACKGROUND ART

In recent years, as a part of thinning to reduce automobile fueleconomy, high strength/thinning of chassis parts have been performed.Due to such thinning, there is an increasing trend in Europe and theAmericas to apply hot-rolled coated steel sheets with improved rustprevention properties to chassis parts normally located inside vehicles.In general, as a hot-rolled steel sheet for automobile chassis parts,steel in which fine precipitates are formed in a ferrite matrixstructure to improve stretch-flangeability has recently been developed(Patent Document 1), but by utilizing a large amount of precipitates toobtain high strength, it may be difficult to obtain high bakehardenability (BH) because the content of solid solution C and Ndecreases.

Accordingly, a technology for a steel sheet capable of securing a bakehardenability value by forming a low-temperature transformationstructure phase by optimizing not only a precipitation strengtheningeffect but also cooling conditions was developed (Patent Document 2).However, Patent Document 2 includes an operation of deforming the steelsheet before secondary cooling after rolling—primary cooling—aircooling, so that it is inevitable to introduce additional facilitiessuch as a temper rolling mill in a ROT section when applied at an actualsite, and there is a problem in that productivity may be lowered due toinferior distribution properties due to the deformation operation. Inaddition, before hot-dip zinc plating after hot rolling, a hard phaseand a dislocation annealing phenomenon are accompanied in an heatingprocess in a range of 450 to 480° C., such that it may be difficult tosecure a sufficient fraction of a Shear Texture in a structure.

Meanwhile, up to now, the issue of securing high bake hardenability incoated steel sheet products has been mainly limited to cold-rolledproducts, and can be divided into two sub-product groups. First, inmanufacturing steel whose tensile strength is 590 MPa or lower, which ismainly applied to exterior panels of automobiles, temper rolling afterplating is applied as an additional technology to improve the bakehardenability (Patent Document 3). However, since the strength of thematerial is fundamentally very low, and the ferrite fraction in thestructure is overwhelmingly high, an effect of increasing dislocationdensity due to physical deformation of the material is not high, precisecontrol during temper rolling does not have a significant effect onimproving bake hardenability. Another product group is a high-strengthcold-rolled steel material that is applied to automobile bodies, and thetechnology related thereto is to secure a low-temperature transformationstructure of an appropriate fraction by controlling a cooling patternafter heating to an austenite transformation temperature after plating,such that it is to improve the bake hardenability through additionalintroduction of dislocation density (Patent Document 4).

However, the prior technologies described above as in Patent Documents 3and 4, that is, content control technologies in ppm units forsolid-solution atoms, are significantly less important in a compositestructure-based hot-rolled coated steel sheet, and in addition, intechnologies for improving bake hardenability by an additionalheat-treatment after plating, there is a need to establish an additionalprocess to be suitable to be applied to a steel sheet having a thicknessof 1 to 5 mm, such that there is a disadvantage of lowering economicefficiency.

Prior Art

(Patent Document 1) Korean Registered Patent Publication No. 10-1203018

(Patent Document 2) Korean Registered Patent Publication No. 10-1657797

(Patent Document 3) Korean Registered Patent Publication No. 10-1676137

(Patent Document 4) Korean Registered Patent Publication No. 10-0691515

DISCLOSURE Technical Problem

An aspect of the present disclosure is to provide a hot rolled coatedsteel sheet having high strength, high formability, and excellent bakehardenability, and a method of manufacturing the same.

Technical Solution

According to an aspect of the present disclosure, a hot rolled coatedsteel sheet having high strength, high formability, and excellent bakehardenability, the hot rolled coated steel sheet includes, in wt %:

C: 0.05 to 0.14%, Si: 0.1 to 1.0%, Mn: 1.0 to 2.0%, P: 0.001 to 0.05%,S: 0.001 to 0.01%, Al: 0.01 to 0.1%, Cr: 0.005 to 1.0%, Ti: 0.005 to0.13%, Nb: 0.005 to 0.03%, N: 0.001 to 0.01%, Fe residues, and otherinevitable impurities; a mixed structure of ferrite and bainite as amain phase; and as a remaining structure, one or more selected from agroup consisting of martensite, austenite, and phase martensite (MA),

wherein a fraction of the ferrite and bainite is 95 to 99 area % andEquation 1 is satisfied,FCO_({110}<112>)+FCO_({112}<111>)≥10   [Equation 1]

where, FCO_({110}<112>) and FCO_({112}<111>), each representing an areafraction occupied by a structure having ac crystal orientation of{110}<112> and {112}<111>.

According to another aspect of the present disclosure, a method ofmanufacturing a hot rolled coated steel sheet having high strength, highformability, and excellent bake hardenability may be provided, themethod comprising operations of:

reheating a steel slab including in wt %, C: 0.05 to 0.14%, Si: 0.1 to1.0%, Mn: 1.0 to 2.0%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, Al: 0.01 to0.1%, Cr: 0.005 to 1.0%, Ti: 0.005 to 0.13%, Nb: 0.005 to 0.03%, N:0.001 to 0.01%, Fe residues, and other inevitable impurities;

hot rolling the reheated steel slab at a temperature of Ar3 or higher to1000° C. to obtain a hot-rolled steel sheet;

primary cooling the hot-rolled steel sheet to a temperature of 550 to750° C.;

ultra-slow cooling the primary-cooled hot-rolled steel sheet to satisfythe following Equation 2;

secondary cooling the ultra-slow cooled hot-rolled steel sheet to atemperature of 300 to 500° C. and then winding the same;

charging the wound hot-rolled steel sheet into a heating table of 350 to550° C. and heating, and then extracting the same;

introducing the heated hot-rolled steel sheet into a hot-dip platingbath of 450 to 550° C. to form a plating layer on a surface of thehot-rolled steel sheet; and

temper rolling the hot-rolled steel sheet on which the plating layer isformed,|T−T _(R)|≤2(T_(R)=241+109[C]+16.9[Mn]+22.7[Cr]−11.1[Si]−5.4[Al]−0.87Temp+0.00068Temp²)  [Equation 2]where T is an actual ultra-slow cooling time, T_(R) is a theoreticalultra-slow cooling time, Temp is an intermediate temperature duringultra-slow cooling, and [C], [Mn], [Cr], [Si], and [Al] mean a contentof each alloy element,15≤(1000−TH)×El_(SPM)≤250   [Equation 3]

where T_(H) is an average temperature of a charging temperature of aheating table of the hot-rolled steel sheet before being charged into ahot-dip plating bath and an extraction temperature of the same, andEl_(SPM) is a difference in length of the hot-rolled steel sheet beforetemper rolling and immediately after temper rolling.

Advantageous Effects

According to an aspect of the present disclosure, it is possible toprovide a hot-rolled coated steel sheet having tensile strength of 780MPa or more and an elongation rate of 10% or more and simultaneouslyhaving excellent bake hardenability, and a method of manufacturing thesame.

DESCRIPTION OF DRAWINGS

FIG. 1 shows EBSD analysis results of Inventive Example 4 andComparative Example 19, (a) shows the EBSD analysis results of InventiveExample 4 and (b)shows the EBSD analysis results of Comparative Example19.

FIG. 2 is a graph showing values of a yield ratio (YR)×bakehardenability (BH) and ductility(El)×stretch-flangeability (HER) ofInventive Examples 1 to 10 and Comparative Examples 1 to 20.

BEST MODE FOR INVENTION

Hereinafter, a hot-rolled coated steel sheet according to an embodimentof the present disclosure will be described.

First, an alloy composition of the hot-rolled coated steel sheet of thepresent disclosure will be described first. A unit of the alloycomposition described below is by weight unless otherwise specified.

Carbon (C): 0.05 to 0.14%

Carbon (C) is the most economical and effective element for reinforcingsteel. If an amount thereof increases, a fraction of low-temperaturetransformation phases such as bainite and martensite increases incomposite structure steel, thereby increasing the tensile strength anddislocation density in the structure. If the content of C is lower than0.05%, it is difficult to easily form a low-temperature transformationphase during cooling after hot rolling, and if the content of C exceeds0.14%, the strength is excessively increased, and there is a problem inthat weldability, formability and toughness are lowered. Therefore, itis preferable that the content of C has a range of 0.05 to 0.14%. Alower limit of the content of C is more preferably 0.06%, and even morepreferably 0.065%. An upper limit of the content of C is more preferably0.13%, even more preferably 0.12%, and most preferably 0.11%.

Silicon (Si): 0.1 to 1.0%

Silicon (Si) deoxidizes molten steel and having a solid solutionstrengthening effect, and silicon (Si) is a ferrite stabilizing element,and has an effect of promoting ferrite transformation during coolingafter hot rolling, and thus silicon (Si) is an effective element forincreasing the ferrite fraction constituting a matrix of the compositestructure steel. If the content of Si is lower than 0.1%, a ferritestabilizing effect is small, and thus it is difficult to form the matrixstructure into a ferrite structure, such that it is difficult to securean elongation rate. If the content of Si exceeds 1.0%, ferritetransformation is excessively promoted, and sufficient dislocation maynot be secured due to a decrease in the fraction of low-temperaturetransformation structure in the structure, and a red scale by Si may beformed on a surface of a steel sheet, so that a surface quality of thesteel sheet is not only very deteriorated, but also weldability arelowered. Therefore, it is preferable that the content of C has a rangeof 0.1 to 1.0%. A lower limit of the content of C is more preferably0.9%, and even more preferably 0.25%. An upper limit of the content ofSi is more preferably 0.9%, even more preferably 0.8%, and mostpreferably 0.7%.

Manganese (Mn): 1.0% to 2.0%

Manganese (Mn) is an effective element for solid solution strengtheningof steel together with the Si. Manganese (Mn) increases hardenability ofsteel to facilitate formation of bainite or martensite during coolingafter hot rolling. However, if the content of Mn is lower than 1.0%, theabove-described effect may not be sufficiently obtained, and if thecontent of Mn exceeds 2.0%, it is difficult to secure an appropriatefraction of ferrite by excessively delaying ferrite transformation, andwhen casing slabs in a casting process, a segregation portion is greatlydeveloped in a thickness central portion, which causesstretch-flangeability to be degraded. Thus, it is preferable that thecontent of Mn has a range of 1.0 to 2.0%. A lower limit of the contentof Mn is more preferably 1.1%, even more preferably 1.2%, and mostpreferably 1.3%. An upper limit of the content of Mn is more preferably1.9%, even more preferably 1.8%, and most preferably 1.7%.

Phosphorus (P): 0.001 to 0.05%

Phosphorus (P) is an impurity present in steel, and if the content of Pexceeds 0.05%, ductility is reduced by micro-segregation and impactproperties of steel are reduced. Meanwhile, in order to control the P to0.001% or lower, it takes a lot of time during a steelmaking operation,so that productivity is greatly reduced. Therefore, it is preferablethat the content of P has a range of 0.001 to 0.05%. The P content ismore preferably 0.001 to 0.04%, even more preferably 0.001 to 0.03%, andmost preferably 0.001 to 0.02%.

Sulfur (S): 0.001% to 0.01%

Sulfur (S) is an impurity present in steel, and if the content of Sexceeds 0.01%, S combines with Mn, or the like to form a non-metallicinclusion, and thus there is a problem that toughness of the steel isgreatly reduced. Meanwhile, in order to control the content of S to belower than 0.001%, a lot of time may be consumed during a steelmakingoperation, resulting in a decrease in productivity. Therefore, it ispreferable that the content of S has a range of 0.001 to 0.01%.Thecontent of S is more preferably 0.001 to 0.007%, even more preferably0.001 to 0.005%, and most preferably 0.001 to 0.003%.

Aluminum (Al): 0.01% to 0.9%

Aluminum (Al) is a component mainly added for deoxidation and ispreferably included in an amount of 0.01% or more in order to expect asufficient deoxidation effect. However, when the content of Al exceeds0.1%, corner cracks are likely to occur in the slab during continuouscasting as AlN is formed by bonding with nitrogen, and defects due toformation of inclusions are likely to occur, which are disadvantages.Therefore, it is preferable that the content of Al has a range of 0.01to 0.1%. A lower limit of the content of Al is more preferably 0.011%,even more preferably 0.013%, and most preferably 0.015%. An upper limitof the content of Al is more preferably 0.08%, even more preferably0.06%, and most preferably 0.05%.

Chromium (Cr): 0.005% to 1.0%

Chromium (Cr) serves to solid strengthen steel, delays ferrite phasetransformation during cooling to facilitate the formation of alow-temperature transformation structure, like Mn. In order tosufficiently obtain the above-described effect, it is preferable that Mnis included in an amount of 0.005% or more. However, if the content ofCr exceeds 1.0%, the ferrite transformation is excessively delayed, sothat a fraction of low-temperature transformation structures such asbainite or martensite phases increases more than necessary, and thus anelongation rate is rapidly deteriorated. Therefore, it is preferablethat the content of Cr has a range of 0.005 to 1.0%. A lower limit ofthe content of Cr is more preferably 0.05%, even more preferably 0.1%,and most preferably 0.2%. An upper limit of the Cr content is morepreferably 0.9%, even more preferably 0.85%, and most preferably 0.8%.

Titanium (Ti): 0.005 to 0.13%

Titanium (Ti) has a representative precipitation strengthening effect,along with Nb, and forms coarse TiN precipitates in steel with strongaffinity with N. The TiN serves to suppress growth of crystal grainsduring a heating process for hot rolling. Meanwhile, Ti remaining afterreacting with N is dissolved in the steel to form TiC precipitates bybonding with C, and the TiC serves to improve the strength of the steel.In order to obtain the above-described effect, it is preferable toinclude Ti in a content of 0.005% or more. However, if the content of Tiexceeds 0.13%, since TiN or TiC precipitates are excessively formed, thefraction of solid-solution atoms such as C and M among steels requiredto obtain high bake hardenability may be rapidly reduced. In addition,due to coarsening of the TiN precipitate, stretch-flangeability maydecrease. Therefore, it is preferable that the content of Ti have arange of 0.005 to 0.13%. A lower limit of the content of Ti is morepreferably 0.01%, even more preferably 0.03%, and most preferably 0.05%.An upper limit of the Ti content is more preferably 0.125%, even morepreferably 0.12%, and most preferably 0.115%.

Niobium (Nb): 0.005 to 0.03%

Niobium (Nb) is a representative precipitation strengthening elementalong with Ti, and serves to improve the strength and impact toughnessof steel by miniaturizing crystal grains through retardation ofrecrystallization by precipitation during hot rolling. In order tosufficiently obtain the above-described effect, it is preferable toinclude the content of Nb to be 0.005% or higher. However, if thecontent of Nb exceeds 0.03%, an amount of solid-solution carbon in thesteel during hot rolling is rapidly reduced, and elongated crystalgrains are formed due to excessive recrystallization delay, resulting inpoor stretch-flangeability, which are disadvantages. Therefore, it ispreferable that the content of Nb has a range of 0.005 to 0.03%. A lowerlimit of the Nb content is more preferably 0.007%, even more preferably0.009%, and most preferably 0.01%. An upper limit of the Nb content ismore preferably 0.025%, even more preferably 0.02%, and most preferably0.018%.

Nitrogen (N): 0.001% to 0.01%

Nitrogen (N)is a representative solution strengthening element togetherwith the C, and forms a coarse precipitate with titanium (Ti), aluminum(Al), and the like. In general, the solid solution strengthening effectof N is better than that of C, but as an amount of N increases in thesteel, there is a problem in that the toughness is greatly reduced.Thus, an upper limit of the N is preferably limited to 0.01%. Meanwhile,if the content of N is 0.001% or lower, it takes a lot of time during asteelmaking operation, and thus productivity is deteriorated. Therefore,it is preferable that the content of N has a range of 0.001 to 0.01%.The content of N is more preferably 0.001 to 0.009%, even morepreferably 0.001 to 0.008%, and most preferably 0.001 to 0.007%.

A remaining component of the present disclosure is iron (Fe). However,in the general manufacturing process, impurities that are not intendedfrom a raw material or a surrounding environment can be inevitablymixed, and therefore cannot be excluded. Since these impurities can beknown to anyone skilled in the ordinary manufacturing process, they arenot specifically mentioned in the present specification.

It is preferable that the hot-rolled coated steel sheet provided by thepresent disclosure includes a mixed structure of ferrite and bainite asa main phase, and includes one or more selected from a group consistingof martensite, austenite, and phase martensite as a remaining structure.The fraction of ferrite and bainite is preferably 95 to 99 area %. Oneor more selected from the group consisting of martensite, austenite andphase martensite (MA) is preferably 1 to 5 area %. By controlling thefraction of the microstructure as described above, the strength,ductility, yield ratio, stretch-flangeability, and bake hardenabilitytargeted by the present disclosure can be secured. If the fraction offerrite and bainite is lower than 95 area % or the fraction of theremaining structure exceeds 5 area %, since a fraction of a hard phaseis excessively increased, and it is difficult to secure sufficientdislocation density in the microstructure due to an increase in anannealing phenomenon of the hard phases during heating before plating,so the bake hardenability decreases, and not only formability such asductility and stretch-flangeability is deteriorated, but alsoweldability is deteriorated, which are problems. On the other hand, whenthe fractions of ferrite and bainite exceed 99 area % or the fraction ofthe remaining structure is lower than 1 area %, since it is not possibleto secure a sufficient fraction of low-temperature transformationstructure in the microstructure, it is difficult to secure sufficientbake hardenability without excessive physical deformation far exceedingan effective El_(SPM) value, presented in Equation 3, which is describedbelow, due to low dislocation density. Here, the low-temperaturetransformation structure is a hard phase such as bainite, martensite,austenite, and phase martensite (MA).

It is preferable that the fraction of bainite is 3 to 30 area%. When thefraction of bainite is lower than 3% by area, the fraction of ferrite ismaximized or the fraction of martensite is increased, so that it is noteasy to secure sufficient dislocation density in the structure. On theother hand, if the fraction thereof exceeds 30% by area, the fraction ofthe hard phase in the structure increases as a whole, resulting in aproblem that the ductility and stretch-flangeability are deteriorated.The fraction of bainite is more preferably 5 to 30%, and even morepreferably 10 to 30%.

It is preferable that the hot-rolled coated steel sheet of the presentdisclosure satisfies the following Equation 1. Typically, in a rolledstructure, α-fibers having a {110}//RD relationship and γ-fibers havinga {111}//ND relationship are mainly observed in the rolled structure,but{110}<112> and {112}<111> crystal orientation is formed when sheardeformation occurs in the structure. So the shear deformation mayfacilitate the generation of dislocations compared to the normal rolledstructure, and may be a measure by which an increase in dislocationdensity in a structure can be measured. The hot-rolled coated steelsheet of the present disclosure can secure excellent strength, ductilityand bake hardenability by introducing sufficient dislocation bysatisfying the following Equation 1. If the value of the followingEquation 1 is not satisfied, a problem in which it is not easy to securehigh bake hardenability may occur because shear deformation sufficientto introduce sufficient dislocation density in the structure is notapplied. Meanwhile, RD and ND mentioned as described above mean aRolling Direction (RD) and a Normal Direction (ND), respectively.FCO_({110}<112>)+FCO_({112}<111>)≥10   [Equation 1]

where, FCO_({110}<112>) and FCO_({112}<111>) refer to the area fractionof the structures having the {110}<112> and {112}<111> crystalorientations, respectively

The hot-rolled coated steel sheet of the present disclosure satisfyingthe above-described alloy composition, microstructure, and Equation 1has bake hardenability (BH): 30 MPa or more, tensile strength (TS): 780MPa or more, an elongation rate (El): 10% or more, a yield ratio (YR):0.8 or more, and stretch-flangeability: 40% or more, so that excellentmechanical properties can be secured. Meanwhile, the bake hardenabilitymay be in accordance with a Low-BH measurement standard.

Meanwhile, the hot-rolled coated steel sheet provided by the presentdisclosure may be a steel sheet in which a plating layer including oneor more of zinc or aluminum on one side or both sides of a base steelsheet, and the plating layer may include all of those commonly used inthe art.

The hot-rolled coated steel sheet of the present disclosure describedabove may be manufactured by various methods, and the manufacturingmethod thereof is not particularly limited. However, as a preferredexample, it may be manufactured by the following methods.

Hereinafter, a method of manufacturing a hot-rolled coated steel sheetaccording to an embodiment of the present disclosure will be described.

(Reheating a Steel Slab)

A steel slab having the above-described alloy composition is reheated. Areheating temperature of the steel slab is preferably 1180 to 1300° C.

If the reheating temperature of the steel slab is lower than 1180° C.,it becomes difficult to secure the temperature during hot rolling due toinsufficient aging heat of the slab, and it becomes difficult to resolvesegregation generated during continuous casting through diffusion. Inaddition, since precipitates deposited during continuous casting may notbe sufficiently resolved, it may be difficult to obtain a precipitationstrengthening effect in a process after hot rolling. On the other hand,if the reheating temperature of the steel slab exceeds 1300° C.,strength reduction and structural unevenness may be promoted by thegrowth of austenite grains. Therefore, it is preferable that thereheating temperature of the steel slab has a range of 1180 to 1300° C.A lower limit of the reheating temperature of the steel slab is morepreferably 1185° C., even more preferably 1190° C., and most preferably1200° C. An upper limit of the reheating temperature of the steel slabis more preferably 1295° C., even more preferably 1290° C., and mostpreferably 1280° C.

(Hot Rolling)

The reheated steel slab is hot-rolled at a temperature of Ar3 or higher,which is a ferrite phase transformation starting temperature, to obtaina hot-rolled steel sheet. If the hot-rolling temperature is lower thanAr3, it is difficult to secure the structure and properties targeted bythe present disclosure by rolling after ferrite transformation, and ifthe hot-rolling temperature exceeds 1000° C., there is a problem in thatformability is deteriorated due to an increase in scale defects increaseon a surface thereof. Therefore, it is preferable that the hot-rollingtemperature has a range of Ar3 or higher to 1000° C. A lower limit ofthe hot-rolling temperature is more preferably 850° C., even morepreferably 860° C., and most preferably 870° C. An upper limit of thehot-rolling temperature is more preferably 935° C., even more preferably930° C., and most preferably 920° C.

(Primary Cooling)

The hot-rolled steel sheet is primarily cooled to a temperature of 550to 750° C. If the primary cooling stop temperature is lower than 550°C., a microstructure of steel mainly includes a bainite phase, so aferrite phase cannot be obtained as a matrix structure, so that it isdifficult to secure a sufficient elongation rate. On the other hand, ifthe primary cooling stop temperature exceeds 750° C., coarse ferrite andpearlite structures are formed, and the desired strength cannot besecured. Therefore, it is preferable that the primary cooling stoptemperature has a range of 550 to 750° C. A lower limit of the primarycooling stop temperature is more preferably 560° C., even morepreferably 580° C., and most preferably 600° C. An upper limit of theprimary cooling stop temperature is more preferably 740° C., even morepreferably 730° C., and most preferably 720° C.

During the primary cooling, a cooling rate is preferably 20° C./sec orhigher. If the primary cooling rate is lower than 20° C./sec, ferriteand pearlite phase transformation occurs during cooling, so that adesired level of hard phase cannot be secured, and desired strength andbake hardenability cannot be secured. Therefore, the primary coolingrate is preferably 20° C./sec or more. The primary cooling rate is morepreferably 30° C./sec or more, even more preferably 45° C./sec or more,and most preferably 60° C./sec or more. Meanwhile, in the presentdisclosure, the higher the primary cooling rate is, the more preferable,so the upper limit of the primary cooling rate is not particularlylimited, and may be appropriately selected in consideration of a coolingfacility.

(Ultra-Slow Cooling)

The primarily-cooled hot-rolled steel sheet is ultra-slow cooled tosatisfy the following Equation 2. The Equation 2 described above is toobtain a microstructure proposed in the present disclosure, and byoptimizing an intermediate temperature (Temp) and an ultra-slow coolingtime during ultra-slow cooling, an appropriate fraction of solidsolution carbon is present after ferrite transformation in a range thatcan secure strength, ductility, and formability. Thereby, an appropriatefraction of low-temperature transformation structure in steel aftercooling is formed, such that it is intended to introduce a sufficientdislocation to an inferior of the low-temperature transformationstructure and an interface of ferrite and the low-temperaturetransformation phase. If the following Equation 2 is not satisfied, itis not possible to secure an appropriate fraction of low-temperaturetransformation structure in the structure of the hot-rolled steel sheet,and thus it may be difficult to secure sufficient bake hardenabilitywithout excessive physical deformation far exceeding the effectiveEl_(SPM) value shown in Equation 3 described below. Therefore, it ispreferable that |T−T_(R)| in the following Equation 2 is 2 or lower. The|T−T_(R)| is more preferably 1.95 or lower, even more preferably 1.925or lower, and most preferably 1.9 or lower. Meanwhile, in the followingEquation 2, T_(R) (a theoretical ultra-slow cooling time) is anultra-slow cooling time to obtain an optimal microstructure fractiontargeted in the present disclosure, and Temp (an intermediatetemperature during ultra-slow cooling) is an intermediate temperature ofan ultra-slow cooling start temperature and end temperature.|T−T _(R)|≤2(T_(R)=241+109[C]+16.9[Mn]+22.7[Cr]−11.1[Si]−5.4[Al]−0.87Temp+0.00068Temp²)  [Equation 2]

where, T is an actual ultra-slow cooling time, T_(R) is a theoreticalultra-slow cooling time, and Temp is an intermediate temperature duringextreme-slow cooling, and [C], [Mn], [Cr], [Si], and [Al] mean a contentof each alloy element

It is preferable that an ultra-slow cooling rate is 2.0° C./sec or lowerduring the ultra-slow cooling. If the ultra-slow cooling rate exceeds2.0° C./sec, there is a disadvantage of causing material deviationbecause a phase transformation behavior of a lengthwise electric fieldof the rolled coil is not uniform. Therefore, the ultra-slow coolingrate is preferably 2.0° C./sec or lower. The ultra-slow cooling rate ismore preferably 1.9° C./sec or lower, even more preferably 1.75° C./secor lower, and most preferably 1.5° C./sec or lower.

An ultra-slow cooling maintaining time during the ultra-slow cooling ispreferably 10 seconds or lower (excluding 0 seconds). If the ultra-slowcooling maintaining time exceeds 10 seconds, the ferrite fractionbecomes also excessively high such that it difficult to secure thedesired strength and bake hardenability. Therefore, the ultra-slowcooling maintaining time is preferably 10 seconds or lower. Theultra-slow cooling maintaining time is more preferably 9.7 seconds orlower, even more preferably 9.5 seconds or lower, and most preferably 9seconds or lower.

An intermediate temperature (Temp) during the ultra-slow cooling ispreferably 545 to 745° C. If the intermediate temperature duringultra-slow cooling is lower than 545° C., the microstructure in thesteel mainly includes a bainite phase, and thus, a ferrite phase cannotbe obtained as a matrix structure, such that it maybe difficult tosecure a sufficient elongation rate. On the other hand, if theintermediate temperature exceeds 745° C., coarse ferrite and pearlitestructure may be formed, such that desired strength may not be secured.That is, the intermediate temperature (Temp) during the ultra-slowcooling is preferably 545 to 745° C. A lower limit of the intermediatetemperature during the ultra-slow cooling is more preferably 550° C.,even more preferably 555° C., and most preferably 560° C. An upper limitof the intermediate temperature during the ultra-slow cooling is morepreferably 740° C., even more preferably 735° C., and most preferably730° C.

(Secondary Cooling)

The ultra-slowly cooled hot-rolled steel sheet is wound after secondarycooling to a temperature of 300 to 500° C. If the secondary cooling stoptemperature is lower than 300° C., a fraction of hard phases such asmartensite, austenite, and phase martensite (MA) is excessivelyincreased, and if the secondary cooling stop temperature exceeds 500°C., a sufficient low-temperature transformation structure fractionincluding bainite may not be secured, such that it is difficult tosecure sufficient bake hardenability without excessive physicaldeformation far exceeding the effective El_(SPM) value presented inEquation 3 described below. Therefore, it is preferable that thesecondary cooling stop temperature has a range of 300 to 500° C. A lowerlimit of the secondary cooling stop temperature is more preferably 310°C., even more preferably 320° C., and most preferably 330° C. An upperlimit of the secondary cooling stop temperature is more preferably 495°C., even more preferably 490° C., and most preferably 485° C.

During the secondary cooling, the cooling rate is preferably 20° C./secor more. If the secondary cooling rate is lower than 20° C./sec, thereis a disadvantage in that it is difficult to secure the strength andbake hardenability suggested by the present disclosure due to anincrease in the ferrite fraction. Therefore, the secondary cooling rateis preferably 20° C./sec or higher. The secondary cooling rate is morepreferably 30° C./sec or more, even more preferably 40° C./sec or more,and most preferably 50° C./sec or more. Meanwhile, in the presentdisclosure, the higher the secondary cooling rate is, the mostpreferable, an upper limit of the secondary cooling rate is not limited,and may be appropriately selected in consideration of a coolingfacility.

(Correcting, Picking)

In the present disclosure, after the winding, an operation of picklingthe wound hot-rolled steel sheet is further included. The pickling isfor removing a scale on a surface of the steel sheet. It is preferablethat the pickling is performed at 200° C. or lower. If the picklingtemperature exceeds 200° C., there is a disadvantage that surfaceroughness of the steel sheet is deteriorated due to excessive pickling.In the present disclosure, a lower limit of the pickling temperature isnot particularly limited, and for example, the lower limit of thepickling temperature may be room temperature. Meanwhile, the steel sheetmay be cooled using natural cooling such as air cooling until thepickling process after the winding. In the present disclosure, beforethe pickling, an operation of correcting the shape of the woundhot-rolled steel sheet may be further included. After the windingprocess, waves may occur at an edge portion of the steel sheet.Therefore, in the present disclosure, it is possible to improve thequality and a yield rate of the steel sheet by correcting the shape ofwaves.

(Heating)

The wound hot-rolled steel sheet is charged to a heating table at 350 to550° C., heated, and then extracted. Control of the heating temperaturecontrol is to improve wettability with a plating solution during asubsequent plating process. If the heating temperature is lower than350° C., there may be a problem in that sufficient wettability is notsecured, so plating properties are deteriorated. If the heatingtemperature exceeds 550° C., a significant amount of dislocations formedin the steel sheet is lost, and it maybe difficult to secure sufficientbake hardenability even if additional dislocations are introduced byphysical deformation to the steel sheet through temper rolling, or thelike. Therefore, the heating temperature is preferably 350 to 550° C. Alower limit of the heating temperature is more preferably 360° C., evenmore preferably 370° C., and most preferably 380° C. An upper limit ofthe heating temperature is more preferably 540° C., even more preferably520° C., and most preferably 500° C.

(Plating, Temper Rolling)

The heated hot-rolled steel sheet is introduced into a hot-dip platingbath of 450 to 550° C. to form a plating layer on a surface of thehot-rolled steel sheet. If the plating temperature is lower than 450°C., since sufficient wettability may not be secured, a problem in thatplating properties are deteriorated, and an elongation rate may alsodecrease. On the other hand, if the plating temperature exceeds 550° C.,a significant amount of dislocations formed in the steel sheet is lost,and it may be difficult to secure sufficient bake hardenability even ifadditional dislocations are introduced by physical deformation to thesteel sheet through temper rolling, or the like. The hot-dip platingbath may include one or more of zinc or aluminum.

During the plating, it is preferable that an introduction speed of thehot-dip plating bath is 10 to 60 mpm (m/min). If the introduction speedof the hot-dip plating bath of the hot-rolled steel sheet is lower than10 mpm, there is a disadvantage such as deterioration of surface qualitydue to overpickling, and if it exceeds 60 mpm, red scale of the surfaceremains due to micropickling, there maybe a disadvantage such asnon-plating. Therefore, it is preferable that the introduction speed ofthe hot-dip plating bath of the hot-rolled steel sheet is 10 to 60 mpm.A lower limit of the hot-dip plating bath of the hot-rolled steel sheetis more preferably 15 mpm, even more preferably 17 mpm, and mostpreferably 20 mpm. An upper limit of the introduction speed of thehot-dip plating bath of the hot-rolled steel sheet is more preferably 58mpm, even more preferably 57 mpm, and most preferably 55 mpm.

Thereafter, temper rolling (SPM) is performed on the hot-rolled steelsheet on which the plating layer is formed. The temper rolling is forintroducing additional dislocations to the hot-rolled steel sheet, andthereby, it is possible to improve the bake hardenability.

Meanwhile, in the present disclosure, it is preferable to satisfy thefollowing Equation 3 in the heating operation and temper rollingoperation described above.

If (1000−T_(H))×El_(SPM) in the following Equation 3 is lower than 15,there is a disadvantage that sufficient dislocations in the structurecannot be secured due to a high heat treatment temperature or a low SPMelongation rate, so that the level of bake hardening ability presentedin the present disclosure cannot be satisfied. If it exceeds 250, thereare disadvantages in that plating quality is deteriorated or such asexcessive strength and insufficient ductility due to excessive SPMworks. Therefore, the (1000−T_(H))×El_(SPM) preferably has a range of 15to 250. A lower limit of (1000−T_(H))×El_(SPM) is more preferably 16,even more preferably 18, and most preferably 20. An upper limit of(1−T_(H))×E_(lSPM) is more preferably 245, even more preferably 240, andmost preferably 230.15≤(1000−T _(H))×El_(SPM)≤250   [Equation 3]

T_(H) is an average temperature of a charging temperature of a heatingtable of a hot-rolled steel sheet before being charged to a hot-dipplating bath and an extraction temperature, and El_(SPM) is thedifference in length of the hot-rolled steel sheet before temper rollingand immediately after temper rolling

Meanwhile, the El_(SPM) is preferably 0.03 to 0.5%.

If the El_(SPM) is lower than 0.03%, the introduction of additionaldislocations is not sufficient, and if El_(SPM) exceeds 0.5%, it maycause deterioration in formability due to a decrease in ductility and anexcessive increase in yield strength. Therefore, the El_(SPM) ispreferably 0.03 to 0.5%. A lower limit of the El_(SPM) is morepreferably 0.04%, even more preferably 0.05%, and most preferably 0.07%.An upper limit of the El_(SPM) is more preferably 0.4%, even morepreferably 0.35%, and most preferably 0.3%.

Mode for Invention

Hereinafter, the present disclosure will be described in more detailthrough examples. However, it is necessary to note that the followingexamples are only intended to illustrate the present disclosure in moredetail and are not intended to limit the scope of the presentdisclosure. This is because the scope of the present disclosure isdetermined by matters described in the claims and able to be reasonablyinferred therefrom.

EXAMPLE

After a steel slab having an alloy composition shown in Table 1 below isprepared, is reheated to 1250° C., and is hot rolled under conditions ofTable 2 to obtain a hot-rolled steel sheet having a 3.5 mm thickness,and then primary cooling, ultra-slow cooling, and secondary cooling wereperformed. In this case, a primary cooling rate was 80° C./sec, and asecondary cooling rate was 70° C./sec. Thereafter, the hot-rolled steelsheet was corrected and pickled, and then plated and temper-rolled underconditions of Table 3 below. After a microstructure and mechanicalproperties on the hot-rolled steel sheet prepared as described above,was measured, results thereof were shown in Tables 3 and 4 below. Inthis case, the microstructure was measured by photographing the steelsheet with a SEM of 3000 magnification, and then an area fraction ofeach phase was calculated using an image analyzer. In particular, thearea fraction of a MA phase in the steel was measured using an opticalmicroscope and a SEM at the same time after etching by a LePera etchingmethod. In addition, in order to analyze a crystal orientation in themicrostructure, an EBSD analysis for an area from 80 μm to 180 μm in athickness direction of the steel sheet, and area to 50 μm in a rollingdirection thereof, that is, for an area of 100 μm×50 μm was performedbased on an interface between a plating layer and a base steel sheet, tomeasure a fraction of structures having {110}<112> and {112}<111>orientations. In addition, mechanical properties were measured bypreparing a DIN standard C-direction specimen for the each hot-rolledsteel sheet, and then performing a tensile test at room temperature at astrain rate of 10 mm/min. Bake hardenability (BH) was measured bypreparing a DIN standard L-direction standard, performing 2%deformation, and then performing heat-treatment the 2% deformed specimenin an oil bath at 170° C. for 20 minutes and air cooling at roomtemperature, and then measuring a low-yield value, as a difference valuethereof. Stretch-flangeability was evaluated based on JFST 1001 to 1996standards.

Alloy composition (weight %) ST No. C Si Mn P S Cr Ti Nb Al N ST 1 0.050.3 1.6 0.01 0.003 0.6 0.09 0.014 0.04 0.003 ST 2 0.07 0.4 1.7 0.010.003 0.4 0.08 0.015 0.05 0.004 ST 3 0.06 0.2 1.7 0.01 0.003 0.5 0.040.014 0.03 0.003 ST 4 0.06 0.3 1.2 0.01 0.003 0.8 0.11 0.015 0.03 0.004ST 5 0.07 0.2 1.2 0.01 0.003 0.7 0.05 0.014 0.04 0.003 ST 6 0.08 0.5 1.50.01 0.003 0.6 0.09 0.014 0.05 0.003 ST 7 0.09 0.7 1.7 0.01 0.003 0.70.04 0.015 0.05 0.003 ST 8 0.09 0.7 1.6 0.01 0.003 0.7 0.08 0.014 0.040.003 ST 9 0.07 0.9 1.9 0.01 0.003 0.6 0.09 0.014 0.04 0.004 ST 10 0.110.9 1.4 0.01 0.003 0.8 0.11 0.015 0.03 0.004 ST 11 0.04 0.8 1.4 0.010.003 0.7 0.08 0.015 0.04 0.003 ST 12 0.17 0.4 1.7 0.01 0.003 0.6 0.110.014 0.04 0.003 ST 13 0.08 0.01 1.5 0.01 0.003 0.7 0.12 0.014 0.050.004 ST 14 0.08 1.2 1.5 0.01 0.003 0.6 0.04 0.015 0.05 0.004 ST 15 0.070.5 0.7 0.01 0.003 0.7 0.11 0.014 0.04 0.003 ST 16 0.08 0.2 2.9 0.010.003 0.4 0.04 0.014 0.04 0.004 ST 17 0.07 0.3 1.6 0.01 0.003 0.004 0.080.014 0.04 0.004 ST 18 0.07 0.8 1.5 0.01 0.003 1.2 0.09 0.014 0.05 0.004ST 19 0.07 0.5 1.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 20 0.07 0.51.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 21 0.07 0.5 1.7 0.01 0.0030.6 0.07 0.014 0.03 0.003 ST 22 0.07 0.5 1.7 0.01 0.003 0.6 0.07 0.0140.03 0.003 ST 23 0.07 0.5 1.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 240.07 0.5 1.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 25 0.07 0.5 1.70.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 26 0.07 0.5 1.7 0.01 0.003 0.60.07 0.014 0.03 0.003 ST 27 0.07 0.5 1.7 0.01 0.003 0.6 0.07 0.014 0.030.003 ST 28 0.07 0.5 1.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 29 0.070.5 1.7 0.01 0.003 0.6 0.07 0.014 0.03 0.003 ST 30 0.07 0.5 1.7 0.010.003 0.6 0.07 0.014 0.03 0.003 *ST: Steel Type

Intermediate temperature Actual during ultra- Secondary TheoreticalFinish Primary ultra-slow slow cooling ultra-slow rolling cooling endcooling cooling end cooling temperature temperature (Temp) time (T)temperature time (T_(R)) Equation Division ST No. (° C.) (° C.) (° C.)(seconds) (° C.) (seconds) 2 IE 1 ST 1  900 640 635 6 450 5.3 0.7 IE 2ST 2  880 600 595 6 470 4.8 1.2 IE 3 ST 3  880 640 635 8 480 7.0 1.0 IE4 ST 4  900 620 615 6 470 4.6 1.4 IE 5 ST 5  890 640 635 6 480 4.1 1.9IE 6 ST 6  860 640 635 6 480 4.6 1.4 IE 7 ST 7  900 600 595 9 480 10.51.5 IE 8 ST 8  900 605 595 8 470 8.8 0.8 IE 9 ST 9  860 640 635 7 4505.9 1.1 IE 10 ST10 870 600 595 8 480 7.7 0.3 CE 1 ST 11 880 640 — 0 470−2.4 2.4 CE 2 ST 12 880 640 635 10 470 19.0 9.0 CE 3 ST 13 880 640 63510 470 12.3 2.3 CE 4 ST 14 880 640 — 0 470 −3.2 3.2 CE 5 ST 15 880 640 —0 470 −7.7 7.7 CE 6 ST 16 880 640 635 10 470 27.1 17.1 CE 7 ST 17 880640 — 0 470 −6.0 6.0 CE 8 ST 18 880 640 635 10 470 13.8 3.8 CE 9 ST 191040 640 635 7 470 7.0 0.0 CE 10 ST 20 840 640 635 7 470 7.0 0.0 CE 11ST 21 880 785 780 10 470 20.4 10.4 CE 12 ST 22 880 525 520 10 470 16.76.7 CE 13 ST 23 880 640 — 0 470 7.0 7.0 CE 14 ST 24 880 640 635 15 4707.0 8.0 CE 15 ST 25 880 640 635 9 600 7.0 2.0 CE 16 ST 26 880 640 635 9200 7.0 2.0 CE 17 ST 27 880 640 635 9 470 7.0 2.0 CE 18 ST 28 880 640635 9 470 7.0 2.0 CE 19 ST 29 880 640 635 9 470 7.0 2.0 CE 20 ST 30 880640 635 9 470 7.0 2.0 [Equation 2] |T − T_(R)|

TABLE 3 Heating Plating bath Equa- Microstructure temperaturetemperature El_(SPM) tion (area %) Division (T_(H)) (° C.) (T_(F)) (°C.) (%) 3 F B M + A + MA IE 1 500 490 0.09 45 83 15 2 IE 2 520 510 0.1572 78 17 5 IE 3 500 490 0.12 60 75 21 4 IE 4 510 500 0.11 54 75 20 5 IE5 500 490 0.13 65 77 21 2 IE 6 490 480 0.16 82 81 16 3 IE 7 500 490 0.1575 82 16 2 IE 8 500 490 0.15 75 77 19 4 IE 9 500 490 0.11 55 78 20 2 IE10 510 500 0.06 29 72 25 3 CE 1 510 500 0.15 74 91 8 1 CE 2 510 500 0.1574 68 22 10 CE 3 510 500 0.15 74 69 21 10 CE 4 510 500 0.15 74 92 6 2 CE5 510 500 0.15 74 93 6 1 CE 6 510 500 0.15 74 64 24 12 CE 7 510 500 0.1574 95 4 1 CE 8 510 500 0.15 74 69 19 12 CE 9 510 500 0.15 74 71 19 10 CE10 510 500 0.15 74 79 10 11 CE 11 510 500 0.15 74 80 12 3 CE 12 510 5000.15 74 69 21 10 CE 13 510 500 0.15 74 71 20 9 CE 14 510 500 0.15 74 879 4 CE 15 510 500 0.15 74 96 2 2 CE 16 510 500 0.15 74 67 19 14 CE 17420 400 0.15 87 65 21 14 CE 18 600 590 0.15 60 93 6 1 CE 19 510 500 0.015 75 21 4 CE 20 510 500 0.91 446 75 21 4 [Equation 3] (1000 − TH) ×El_(SPM) T_(h) is an average temperature of a charging temperature of aheating table of the hot-rolled steel sheet before being charged into ahot-dip plating bath and an extraction temperature of the same, El_(SPM)is a difference in a length of the hot-rolled coated steel sheet beforetemper rolling and immediately after temper rolling. F: Ferrite, B:Bainite, M: Martensite, A: Austenite, MA: Phase martensite

TABLE 4 Crystal orientation (area %) Equa- BH {110} {112} tion YS TS ElHER (MPa) Division <112> <111> 1 (MPa) (MPa) YR (%) (%) BH IE 1 5 6 11657 782 0.84 15 59 40 IE 2 7 6 13 674 793 0.85 13 65 50 IE 3 10 6 16 703790 0.89 14 52 62 IE 4 9 8 17 696 782 0.89 14 55 61 IE 5 7 7 14 706 8020.88 14 57 54 IE 6 6 8 14 687 799 0.86 15 55 44 IE 7 6 5 11 705 810 0.8715 60 57 IE 8 6 7 13 676 805 0.84 15 57 59 IE 9 8 5 13 729 819 0.89 1449 54 IE 10 9 10 19 774 850 0.91 13 48 61 CE 1 5 3 8 473 520 0.91 13 8821 CE 2 4 4 8 574 809 0.71 9 40 22 CE 3 4 3 7 554 770 0.72 9 55 19 CE 43 3 6 664 730 0.91 16 53 12 CE 5 3 5 8 653 710 0.92 17 60 15 CE 6 3 4 7600 870 0.69 10 41 18 CE 7 2 2 4 644 700 0.92 18 45 9 CE 8 1 3 4 650 8900.73 10 39 7 CE 9 7 7 14 503 752 0.67 13 31 45 CE 10 7 6 13 698 812 0.868 49 34 CE 11 3 2 5 671 818 0.82 9 55 19 CE 12 4 3 7 748 880 0.85 9 5125 CE 13 3 3 6 689 840 0.82 10 43 18 CE 14 3 2 5 654 808 0.81 15 46 11CE 15 2 2 4 692 778 0.89 15 49 16 CE 16 7 7 14 632 890 0.71 7 40 49 CE17 7 6 13 616 893 0.69 7 39 37 CE 18 2 0 2 703 772 0.91 15 45 2 CE 19 31 4 643 794 0.81 15 43 12 CE 20 16 15 31 856 911 0.94 7 36 83 [Equation1] FCO{110} <112> + FCO {112} <111> YS: Yield strength, TS: Tensilestrength, YR: Yield rate, El: Elongation rate, HR:Stretch-flangeability, BH: Bake hardenability *IE: Inventive Example CE:Comparative Example

As can be seen from the above Tables 1 to 4, in the case of InventiveExamples 1 to 10 satisfying an alloy composition, a microstructure,manufacturing conditions, and Relational equations 1 to 3 proposed bythe present disclosure, it can be seen that excellent mechanicalproperties are secured, such as bake hardenability (BH): 30 MPa or more,tensile strength (TS): 780 MPa or more, an elongation rate (El): 10% ormore, a yield ratio (YR): 0.8 or more, and stretch-flangeability: 40% ormore.

Comparative Examples 1 to 8 show cases in which the alloy compositionproposed by the present disclosure is not satisfied, as a content of C,Si, Mn, and Cr, which greatly contributes to formation of ferrite and alow-temperature transformation structure, was deviated, a microstructurefraction or the following equation 1 proposed by the present disclosurewas not satisfied. To this end, it can be seen that the presentdisclosure does not secure desired physical properties.

Comparative Examples 9 and 10 show cases in which the alloy compositionproposed by the present disclosure is satisfied, but a finish rollingtemperature does not satisfy the conditions of the present disclosure.In the case of Comparative Example 9, formability is deteriorated due togeneration of excessive red scales, and in the case of ComparativeExample 10, the temperature is controlled to lower than Ar3, it is noteasy to secure stretch-flangeability due to formation of an elongatedstructure by rolling during ferrite transformation.

Comparative Examples 11 and 12 show cases in which the alloy compositionproposed by the present disclosure is satisfied, but a first coolingstop temperature does not satisfy the conditions of the presentdisclosure, and it can be seen that it is difficult to secure themicrostructure fraction proposed by the present disclosure, such thatthe mechanical properties to be obtained by the present disclosure arenot secured. In particular, in the case of Comparative Example 11, asthe first stop temperature exceeds 750° C. to form a 5% pearlitestructure, it can be seen the mechanical properties to be obtained bythe present disclosure are not secured.

Comparative Examples 13 and 14 show cases in which the alloy compositionproposed by the present disclosure is satisfied, but an ultra-lowcooling holding time does not satisfy the conditions of the presentdisclosure, and it can be seen that the mechanical properties to beobtained by the present disclosure are not be secured as the Relationalequation 2 is not satisfied.

Comparative Examples 15 and 16 show cases in which the alloy compositionproposed by the present disclosure is satisfied, but a secondary coolingstop temperature does not satisfy the conditions of the presentdisclosure, and it can be seen that it is difficult to secure themicrostructure fraction proposed by the present disclosure, such thatthe mechanical properties to be obtained by the present disclosure arenot secured.

Comparative Example 17 shows a case in which the alloy compositionproposed by the present disclosure is satisfied, but a plating bathtemperature does not satisfy the conditions of the present disclosure,and it can be seen that an elongation rate is at a low level.

Comparative Example 18 shows a case where the alloy composition proposedby the present disclosure is satisfied, but a heating temperature and aplating bath temperature do not satisfy the conditions of the presentdisclosure, and it can be seen that bake hardenability is at a lowlevel.

Comparative Examples 19 and 20 show cases in which the alloy compositionproposed by the present disclosure is satisfied, but El_(SPM) does notsatisfy the conditions of the present disclosure, and it can be seenthat bake hardenability or elongation is at a low level.

FIG. 1 shows EBSD analysis results of Inventive Example 4 andComparative Example 19, in FIG. 1 , (a) shows the EBSD analysis resultsof Inventive Example 4, and (b) shows EBSD analysis results ofComparative Example 19. In the case of Inventive Example 4, a largeamount of structures having the {110}<112> and {112}<111> crystalorientations proposed by the present disclosure were formed, whereas inComparative Example 19, it can be seen that formation of a structurehaving a crystal orientation of the {110}<112> and {112}<111> is notsufficient.

FIG. 2 is a graph showing a value of a yield ratio (YR)×bakehardenability (BH) and elongation (El)×stretch-flangeability (HER) ofInventive Examples 1 to 10 and Comparative Examples 1 to 20. As can beseen from FIG. 2 , in the Inventive Examples of the present disclosure,it can be seen that a yield ratio (YR), bake hardenability (BH),ductility (El), and elongation flangeability (HER) are all superior tothe Comparative Examples.

The invention claimed is:
 1. A hot-rolled coated steel sheet comprising: in wt %, C: 0.05-0.14%, Si: 0.1-1.0%, Mn: 1.0-2.0%, P: 0.001-0.05%, S: 0.001-0.01%, Al: 0.01-0.1%, Cr: 0.005-1.0%, Ti: 0.005-0.13%, Nb: 0.005-0.03%, N: 0.001-0.01%, and a balance of Fe and inevitable impurities, wherein the hot-rolled coated steel sheet comprises a mixed structure of ferrite and bainite as a main phase, and as a remaining structure, one or more selected from a group consisting of martensite, austenite, and phase martensite (MA), wherein a fraction of the ferrite and bainite is 95 to 99 area%, and Equation 1 is satisfied, FCO_({110}<112>)+FCO_({112}<111>)≥10   [Equation 1] where, FCO_({110}<112>) and FCO_({112}<111>), each represents an area fraction occupied by a structure having ac crystal orientation of {110}<112> and {112}<111>.
 2. The hot-rolled coated steel sheet of claim 1, wherein a fraction of bainite is 3 to 30 area%.
 3. The hot-rolled coated steel sheet of claim 1, wherein the hot-rolled coated steel sheet has bake hardenability (BH): 30 MPa or more, tensile strength (TS): 780 MPa or more, an elongation rate (El): 10% or more, a yield ratio (YR): 0.8 or more, and stretch-flangeability: 40% or more.
 4. The hot-rolled coated steel sheet of claim 1, wherein in the hot-rolled coated steel sheet, a plating layer including one or more of zinc or aluminum is formed on one side or both sides of a base steel sheet. 